NiCrCoMoW Age Hardenable Alloy for Creep-Resistant High Temperature Applications, and Methods of Making

ABSTRACT

The invention provides a Ni-based superalloy with good yield stress and ultimate tensile strength and good creep strength (long creep life at high temperature). Methods of making the alloy are also described.

GOVERNMENT RIGHT CLAUSE

This invention was made with Government support under contract 89243318CFE000003 awarded by the U.S. Department of Energy. The Government has certain rights in this invention.

RELATED APPLICATIONS

This application claims the priority benefit of U.S. Provisional patent application Ser. No. 63/326,270 filed Mar. 31, 2022.

INTRODUCTION

Future power plant designs, namely advanced ultra-supercritical (AUSC) and/or supercritical CO₂ (sCO₂) power plants are expected to raise the efficiencies of coal-fired power plants from about 35% to >50% and decrease harmful gas emissions. In order to achieve this goal, the operating temperatures of high-yield stress components in the coal-fired boiler and steam turbine must be increased. Alternatively, next generation gas turbines need alloys for rotors that possess high-yield stress and environmental resistance to crack formation and growth (i.e., for cycling operation) at temperatures up to 800° C. to minimize efficiency losses due to steam cooling to achieve the desired 70% efficiency target. Critical components that can resist deformation at temperatures above 800° C. are becoming essential and life-limiting. These components are often made from advanced powder metallurgy Ni-based superalloys; however, powder techniques limit the size of components and increase manufacturing cost. For this reason, cast/wrought alloys are preferred, particularly in electricity generation.

The cast and wrought Ni-based superalloy HAYNES® 282® (Table 1A and 1B) has been one of the best candidate alloy for high-temperature, creep resistant, components in advanced power plants. The typical microstructure of HAYNES® 282® consists of a face-centered cubic (FCC) matrix with γ′ precipitates embedded within the γ′ matrix grains when heat treated appropriately. The γ′ precipitates are Ni₃(Al, Ti) with ordered FCC Ll₂ crystal structure. Carbides such as MC and grain boundary M₂₃C₆ are also found within the microstructure.

TABLE 1A Composition range for HAYNES ® 282 ® (wt. %) (http://haynesintl.com/docs/default- source/pdfs/new-alloy-brochures/high- temperature-alloys/brochures/282- brochure.pdf?sfvrsn=20; https://www.elgiloy.com/strip-haynes-alloy-282/) Element Ni Cr Co Mo Ti Al C B Fe Mn Si P Cu Nominal Bal. 20 10 8.5 2.1 1.5 0.06 0.005 — — — — — Minimum Bal. 18.5 9.0 8.0 1.90 1.38 0.04 0.003 — — — — — Maximum Bal. 20.5 11.0 9.0 2.30 1.65 0.08 0.010 1.5 0.3 0.15 0.015 0.1

TABLE 1B Composition range for HAYNES ® 282 ® from patent (wt. %) with impurities balanced (https://patents.google.com/patent/US806693882/en) Al + 0.56TI + Mo + Element Ni Cr Co Mo W Ti Al Nb Ta C B Zr 0.29Nb 0.52W Minimum Bal. 17 8 4.0 — 1.50 1.39 — — 0.01 — — 2.2 6.5 Maximum Bal. 22 15 9.1 7.0 2.30 1.65 0.80 1.5 0.2 0.015 0.02 2.9 9.1

HAYNES® 282® has been a favorite candidate for AUSC applications for a number of aspects. Its strength surpasses that of other candidate alloys (such as precipitation strengthened HAYNES® 263® and INCONEL® 740H® and solid solution strengthened HAYNES® 230® and INCONEL® alloy 617). It is less than that of high-strength Waspaloy® or René-41 Ni-based superalloys, however, HAYNES® 282® shows greater fabricability, an important attribute for AUSC applications and the fabrication of large components or tubing. HAYNES® 282® is also weldable. As a result, HAYNES® 282® also found applications in other power generation applications such as gas turbines in commercial aircraft.

Although HAYNES® 282® possesses attractive creep and tension properties, the performance of the alloy starts to deteriorate rapidly past ˜750° C. and the alloy underperforms when compared to more advanced Ni-based superalloys, particularly those manufactured from powder metallurgy techniques. The invention provides an alloy similar to HAYNES® 282® but having properties at 800° C. similar to HAYNES® 282® at 760° C. and at least 30% improvement in creep life. The alloy is easily fabricable using a cast/wrought processing route, weldable and shows good phase stability at temperatures up to at least 800° C.

SUMMARY OF THE INVENTION

In a first aspect, the invention provides a nickel alloy, comprising: 18.5 to 20.5% Cr; 9.0 to 11.0% Co; 4.0 to 5.0% Mo; 3.25 to 4.25% W; 2.3 to 2.7% Ti; 1.5 to 1.9% Al; 0 to 0.1% C; 0 to 0.02% B; up to 1.5% Fe, 0.5% or less of Mn, Si, P, Cu or other transition metals; the balance Ni and impurities.

In another aspect, the invention provides a nickel alloy, comprising: 18.5 to 20.5% Cr; 9.0 to 11.0% Co; greater than 50% Ni and characterizable by one or any combination of the following properties: a time to failure at 760° C. and 259 MPa of at least 4000 h, or at least 5000 h, or at least 6000 h, or at least 7000 h, or in the range of 4000 to 8000 h or in the range of 4500 to 7500 h (this property refers to measurements in creep using the ASTM standard E-139); and a time to failure at 800° C. and 155 MPa of at least 3000 h, or at least 4000 h, or at least 5000 h, or at least 6000 h, or in the range of 3000 to 7000 h or in the range of 4000 to 6500 h (this property refers to measurements in tension using the ASTM standard E-8). Preferably, the alloy comprises γ′ (or γ″ or γ′/γ″) precipitates distributed within a y matrix grain interior.

In a further aspect, the invention provides a nickel alloy, comprising: 18.5 to 20.5% Cr; 9.0 to 11.0% Co; greater than 50% Ni and characterizable by one or any combination of the following properties: a time to failure at 800° C. and 259 MPa of at least 700 h, or at least 800 h, or at least 900 h, or at least 1000 h, or in the range of 700 to 2000 h or in the range of 900 to 1500 h (this property refers to measurements in creep using the ASTM standard E-139); and a time to failure at 840° C. and 155 MPa of at least 1200 h, or at least 1400 h, or at least 1600 h, or at least 1800 h, or in the range of 1200 to 2200 h or in the range of 1600 to 2000 h (this property refers to measurements in tension using the ASTM standard E-8). Preferably, the alloy of claim D comprising γ′ (or γ″ or γ′/γ″) precipitates distributed within a y matrix grain interior.

In yet another aspect, the invention provides a nickel alloy comprising 18.5 to 20.5% Cr; 9.0 to 11.0% Co and characterizable by one, or any combination of the properties described in the specification.

In any of its aspects, the alloy can be further characterized by one or any combination of the following: comprising: 18.5 to 20.5% Cr; 9.0 to 11.0% Co; 4.0 to 5.0% Mo; 3.25 to 4.25% W; 2.3 to 2.7% Ti; 1.5 to 1.9% Al; 0.003 to 0.01% B; 0.0 to 1.5% Fe; 0.04 to 0.08% C; 0 to 0.3% Mn; 0 to 0.15% Si; 0 to 0.15% P; 0 to 0.01% Cu; the balance Ni and impurities; comprising

M₂₃C₆ carbides at grain boundaries in the alloy; comprising γ′ (or γ″ or γ′/γ″) precipitates distributed within a y grain; characterizable by a time to failure at 760° C. and 259 MPa of at least 4000 h, or at least 5000 h, or at least 6000 h, or at least 7000 h, or in the range of 4000 to 8000 h or in the range of 4500 to 7500 h (measurements in creep use the ASTM standard E-139); and/or characterizable by a time to failure at 800° C. and 155 MPa of at least 3000 h, or at least 4000 h, or at least 5000 h, or at least 6,000 h, or in the range of 3000 to 7000 h or in the range of 4000 to 6500 h.

In another aspect, the invention provides a method of heat treating a nickel alloy comprising Cr and Co, comprising the sequential steps of treating the nickel alloy: (1100±25° C. for 0.5 to 5 h)+(1160±25° C. for 1.5 to 10 h)+(1200±25° C. for 2 to 20 h).

In any of its aspects, the alloy can be further characterized by one or any combination of the following: wherein the nickel alloy comprises any of the alloys described herein; further comprising a heat treatment comprising heating at 1165±25° C. for 400±100 sec and aging steps of 1010±20° C. for 0.5 to 10 h and then 788±15° C. for 2 to 15 h; further comprising a heat treatment comprising heating at 1165±25° C. for 400±100 sec and aging steps of 1060±20° C. for 0.5 to 10 h and then 800±15° C. for 2 to 15 h; further comprising a heat treatment comprising heating at 1165±25° C. for 400±100 sec and an aging step of 800±15° C. for 2 to 15 h; wherein, relative to HAYNES® 282® alloy, the treatments increase life of the alloy in creep at 760° C. and 259 MPa by at least 70%, or at least 100%, or in the range of 80 to 150%, 90 to 140% or 80 to 130%; and/or wherein the treatments increase the elongation to failure of the alloy in creep at 780° C. and 207 MPa by at least 40%, or at least 60%, or in the range of 40 to 100%, 50 to 90% or 40 to 80%.

The invention also includes an alloy made by any of the methods described herein. Any aspect of the invention may be further characterized by having the one or any combination of the properties (or within ±30% or ±20% or ±10% or ±5%) of one or any combination of the properties described herein.

Various aspects of the invention are described using the term “comprising;” however, in narrower embodiments, the invention may alternatively be described using the terms “consisting essentially of” or, more narrowly, “consisting of

The invention also comprises articles comprising a component in: steam turbines, coal fired boilers and gas turbines, including compressor blades, turbine blades, turbine discs and spacers, turbine vanes, combustion cans, boiler tubes, boiler pipes, headers, power generation systems using fluids such as supercritical carbon dioxide (e.g. advanced ultra-supercritical power plants), concentrated solar power plants, nuclear power plants, molten salt reactors: casings, valves, heat exchangers and/or recuperators; made of an alloy of the invention.

The Invention includes uses including: Use as heat resistant structural material for components that require high yield stress at room temperature and good creep strength at high temperatures such as articles found in steam turbines, coal fired boilers and gas turbines, including compressor blades, turbine blades, turbine discs and spacers, turbine vanes, combustion cans, boiler tubes, boiler pipes, headers.; use as structural material for components in aero engines (gas fan turbines) or power generation systems using fluids such as supercritical carbon dioxide (e.g., advanced ultra-supercritical power plants), concentrated solar power plants, nuclear power plants, molten salt reactors: turbine blades, casings, valves, heat exchangers and recuperators.

BRIEF DESCRIPTION OF THE FIGURES

FIG. 1 . (a) Thermo-Calc phase predictions for the alloys with a dashed line representing the γ′ fraction at 800° C. in STD and equivalent fraction in alloys O, P and Q. (b) Microstructure of the alloys following 100 h exposure at 800° C.

FIG. 2 . Larson-Miller plots for alloys Q (invention), STD (NETL-processed HAYNES® 282® formulation) and the report from ORNL and (right) a comparison to other commercial Ni-based superalloys: NIMONIC® alloy 263 and INCONEL® alloy 725 types processed at NETL and INCONEL® alloy 740 from http://dx.doi.org/10.1016/j.msea.2013.04.087.

FIG. 3 . Tensile properties (UTS, 0.2% yield stress and elongation to failure) for alloy Q (invention) with SA and SSA heat treatments, for the commercial alloy from the HAYNES® 282® datasheet (Haynes 282) and for the commercial alloy with SSA heat treatment from a Oak Ridge National Laboratory code case (ORNL 2020: https://doi.org10.2172/1649169)

FIG. 4 . Microstructures of (a) the standard alloy (STD) and (b) alloy Q following SA and of alloy Q in (c) the as-rolled condition and (d) after exposure at 1010° C. for 2 h.

FIG. 5 . Microstructures of the standard alloy (STD) and alloy Q following 5,000 h at 800° C.

FIG. 6 . Microstructures of the standard alloy (STD) and alloy Q following 500 h at 900° C.

DETAILED DESCRIPTION OF THE INVENTION

The invention provides a cast/wrought alloy of composition within the range of major elements listed in Table 2. The alloy can be subjected to a computationally optimized homogenization heat treatment.

TABLE 2 Composition range for the invention(wt. %). Element Ni Cr Co Mo W Ti Al C B Fe Mn Si P Cu Nominal Bal. 20 10 4.5 3.85 2.5 1.7 0.06 0.005 — 0.2 0.05 — — Minimum Bal. 18.5 9.0 4.0 3.25 2.3 1.5 0.04 0.003 — — — — — Maximum Bal. 20.5 11.0 5.0 4.25 2.7 1.9 0.08 0.010 1.5 0.3 0.15 0.015 0.1 Note: Bold. underlined text shows concentrations outside of the HAYNES ® 282 ® range.

Compared to nominal HAYNES® 282® (Table 1A), both Ti and Al concentrations were increased to target a γ′ fraction at 900° C. similar to that of HAYNES® 282® at 800° C. Thereby, the level of precipitate strengthening is increased in the invention. Furthermore, part of the Mo content was substituted for W for phase stability considerations.

An alloy named Q with a target chemistry matching that of the nominal from Table 2 was manufactured following melt processing techniques of an industrial nature. The ingot of approximately 8 kg was made at NETL using high-purity industry-grade melt materials. The alloy was vacuum induction melted under Ar partial pressure to about 50° C. above the liquidus temperature predicted using Thermo-Calc. The liquid was then poured in a graphite mold with ceramic wash coat to prevent C pickup and solidified as a cylindrical ingot. The ingot was then homogenized using the NETL approach, which assesses the chemistry in a computational manner so as to minimize heterogeneity. The heat treatment schedule used to create homogeneity throughout the microstructure comprises the following: 1100° C. for 1 h followed by 1160° C. for 3 h followed by 1200° C. for 12 h

This homogenization heat treatment schedule incorporates thermodynamic factors as well as kinetic ones (i.e., time for specific elements to diffuse within the γ′ matrix) to complement microstructural considerations (i.e., spacing of dendrite arms in the solidified microstructure as well as the interdendritic spacing between those arms), and the sequence described encompasses homogeneity within the alloy from at least 1% to 10%.

The ingot was then hot worked at 1165±25° C. using steps of forging followed by hot rolling with reheating between each step to form a 10 mm thick plate. The last reheat thermal cycle (following the last hot rolling step) was used as a solution heat treatment which lasted 400 ±100 seconds.

The plate was then aged following the heat treatment commonly used for commercial HAYNES® 282® alloy, referred to as standard aging (SA), which comprises the following: 1010° C. for 2 h followed by air cooling and then 788° C. for 8 h followed by air cooling.

The first thermal cycle in the heat treatment is used to form M₂₃C₆ carbides along the grain boundaries and control their morphology. The second step allows for the precipitation of fine γ′ precipitates. A second aging heat treatment was considered and referred to as single step aging (SSA) which comprises the following: 800° C. for 4 h followed by air cooling.

Additional alloys were manufactured following the details above with compositions outside the range expressed for the invention in Table 2. These are listed in Table 3 with compositions obtained using x-ray fluorescence and combustion analysis alongside alloy Q. Alloy STD is the nominal HAYNES® 282® composition used for comparison to the invention (alloy Q). In Alloys O and P, the fraction of γ′ precipitates was designed such that at 900° C. it is similar to that of HAYNES® 282® at 800° C. However, a phase stability study revealed the formation of undesirable phases (likely TCP phases) along the grain boundaries of O and P, observed in FIG. 1 b as bright blocky precipitates. The microstructures were obtained on a scanning electron microscope (SEM). Electron dispersive spectroscopy (EDS) further revealed the precipitates to be rich in Mo which led to the formulation of alloy Q in which part of the Mo was substituted with W. This resulted in a compromise between phase stability (Figure 1 b in which no TCP phase was observed after 100 h at 800° C.) and γ′ fraction/solvus temperature (FIG. 1 a ).

TABLE 3 Composition of the various heats from XRF analysis for the main elements, combustion analysis for C and calculated from the melt addition for B (wt. %). Alloy Ni Cr Co Mo W Ti Al C B Fe Mn Si Cu STD 57.9 19.9 9.7 8.5 <0.01 2.1 1.7 0.06 0.005 <0.01 0.11 0.04 <0.003 O 57.0 19.9 9.7 8.4 <0.01 2.5 2.1 0.06 0.005 <0.01 0.19 0.04 <0.003 P 57.1 19.9 9.7 8.5 <0.01 2.5 2.1 0.06 0.005 <0.01 0.20 0.05 <0.003 Q 57.4 19.9 9.8 4.5 3.85 2.5 1.7 0.06 0.005 <0.01 0.21 0.05 0.02

Properties

Mechanical testing was conducted in creep at various temperatures between 740° C. and 840° C. and stresses between 310 MPa and 69 MPa using the ASM standard E-139 and in tension at room temperature, 450° C., 675° C. and 800° C. using the ASM standard E-8.

The creep properties of the alloys are listed in Table 4 and the results are discussed below using FIG. 2 to facilitate the observation of various trends. The invention (alloy Q) is compared to the commercial HAYNES® 282 0 alloy using the NETL-processed standard alloy (STD) and a large databased from Oak Ridge National Laboratory (ORNL 2020 available in the following: (https://doi.org/10.2172/1649169).

TABLE 4 Creep properties of the invention (alloy Q) and the standard HAYNES ® 282 ® formulation (STD). Aging Heat Temperature Stress Time to Failure LMP Elongation to Alloy Treatment (° C.) (MPa) (h) (C = 20)/1000 Failure (%) Q SA 740 310.26 7,172 24.169 11.6 760 258.55 7,406 24.661 11.7 780 206.84 7,339 25.134 15.4 800 155.13 6,231 25.535 — 820 103.42 21,618 26.602 9.2 840/900 68.95 9,865/2,980* 26.587* 10.1* SSA 760 258.55 5,897 24.559 8.6 800 155.13 9,143 25.714 6.9 800 258.55 1,090 24.723 10.6 840 155.13 1,808 25.889 8.5 900 124.10 361 26.463 9.8 STD SA 760 258.55 3,227 24.285 9.7 780 206.84 4,203 24.876 18.2 900 124.10 192 26.142 16.1 *Temperature was changed from 840° C. after 9,865 h to 900° C. Life at 900° C./69 MPa was converted using LMP which resulted in 2,980 h.

A few important points are listed below:

-   -   The life of alloy Q increased by 130% compared to the         NETL-processed commercial alloy at 760° C. and 259 MPa for         similar aging heat treatments.     -   The life of alloy Q increased by 75% compared to the         NETL-processed commercial alloy at 780° C. and 207 MPa for         similar aging heat treatments.     -   The life of alloy Q with the SSA aging increased by 83% compared         to the NETL-processed commercial alloy at 760° C. and 259 MPa.     -   The life of alloy Q with the SSA aging increased by 88% compared         to the NETL-processed commercial alloy at 900° C. and 124 MPa.     -   Alloy Q outperformed candidate Ni-based superalloys for AUSC or         sCO₂ applications INCONEL® alloy 740 (data from         http://dx.doi.org/10.1016/j.msea.2013.04.087), and         NETL-processed NIMONIC® alloy 263 and INCONEL® alloy 725         equivalents (FIG. 2 ).

Results from tension testing are compiled in Table 5. Similar to the description of the creep results, FIG. 3 is provided below to better illustrate the trends. The invention (alloy Q) is compared to the properties reported for HAYNES® 282® in the manufacturer's datasheet http://haynesintl.com/docs/default-source/pdfs/new-alloy-brochures/high-temperature-alloy/brochures/282-brochure.pdf?sfvrsn=20).

TABLE 5 Tensile properties of the invention and commercial alloy HAYNES ® 282 ®. Ultimate 0.2% Aging Tensile Yield Elongation Heat Temperature Strength Stress to Failure Alloy Treatment (° C.) (MPa) (MPa) (%) Q SA 24 1,230 780 17.5 450 1,108 745 18.9 675 1,147 733 16.6 800 815 674 18.9 SSA 24 1,304 873 30.0 450 1,117 843 22.5 675 1,064 803 10.1 800 840 733 4.3 HAYNES ® SA 24 1,147 715 30 282 ® 538 991 649 34 649 1,048 643 31 704 978 649 29 760 856 628 22 816 709 575 28

From FIG. 3 , the ultimate tensile strength (UTS) and 0.2% yield stress (YS) increased from the standard alloy to alloy Q. On the other hand, the elongation to failure decreased. A few important points are listed below:

-   -   The UTS and YS were increased by 7% and 9% at room temperature,         respectively, while the elongation to failure decreased by 42%         (Comparing HAYNES® 282® and Alloy Q SA).     -   The UTS and YS were increased by 8.5% and 14% at 800° C.,         respectively, while the elongation to failure decreased by 28%         (Comparing HAYNES® 282® and Alloy Q SA). Note: The values for         the commercial alloy at 800° C. were determined using linear         interpolations with the values at 760° C. and 816° C.     -   The YS of Alloy Q with the SA heat treatment was higher than         with the SSA heat treatment.

The microstructures of the alloys following SA are compared in FIG. 4 a and b. Fine γ′ precipitates are present within the y grains in both alloys (STD and Q) but hard to resolve on the images. Carbides are decorating the grain boundaries and were identified as M₂₃C₆ carbides using energy dispersive spectroscopy. Alloy Q contained larger γ′ precipitates as well as the fine ones, FIG. 4 b . These were formed during the first step of the SA heat treatment consisting of 1010° C. for 2 h as illustrated in FIG. 4 c and d where only fine γ′ precipitates are present after rolling (likely forming during cooling to room temperature), FIG. 4 c , while the coarse γ′ precipitates appeared after exposing the alloy to 1010° C. for 2 h, FIG. 4 d .

The bimodal-type size distribution of the γ′ precipitates in alloy Q following SA is likely not ideal for long term mechanical performance. An alternative approach is to eliminate the first step of the heat treatment, similar to the SSA, to produce the fine γ′ precipitates. A second approach is to raise the temperature of the first step of the SA, originally designed for grain boundary carbide distribution, to above the γ′ solvus temperature. Heat treatment trials revealed the γ′ solvus temperature to be above 1030° C.

A phase stability study was conducted and selected images for the standard alloy (STD) and alloy Q are compared in FIGS. 5 and 6 . Detrimental phases identified as μ and/or σ (TCP phases) formed in significant amounts in the standard alloy following as early as 1,000 h exposure at 800° C. The low magnification inset of FIG. 5 shows the extent of formation of the TCP phases in the standard alloy. Significantly less precipitation of TCP phases was observed in alloy Q (FIG. 5 ). Similarly, the STD was more prone to the formation of TCP at 900° C. than alloy Q, as shown in FIG. 6 following 500 h exposure. TCP phases are detrimental to Ni-base superalloys, particularly when resistance to creep deformation is required as it involves significant time at the operating temperature.

Summary of Results

The invention provides a material's chemistry range (Table 2) and a homogenization heat treatment formulated for the refined chemistry to improve the mechanical properties (particularly resistance to creep deformation) of HAYNES® 282® alloy past 800° C.

The homogenization heat treatment provides an initial uniform chemistry throughout which avoids elemental segregation that can lead to the formation of undesirable phases over time (e.g., during extended exposure time at temperature in creep).

The life of alloy Q increased by 130% compared to the commercial alloy at 760° C. and 259 MPa. The life of alloy Q increased by 75% compared to the commercial alloy at 780° C. and 207 MPa. The life of alloy Q with the SSA aging increased by 163% compared to the NETL-processed commercial alloy for an applied stress of 259 MPa. For the comparison, alloy Q with SSA aging was tested at 800° C. and the resulting life was converted to equivalent life at 760° C. using the Larson-Miller parameter.

Alloy Q outperformed candidate Ni-based superalloys for AUSC or sCO2 applications Inconel 740, data taken from literature, Nimonic 263 and Inconel 725 equivalents processed at NETL (FIG. 2 ).

Phase stability, as defined by avoiding or delaying the formation of undesirable phases (such as TCP), was improved in alloy Q compared to the commercial formulation.

The alloy showed great workability as cold rolling tests were performed and successful (e.g., no visible cracks on the rolled plate).

The homogenization heat treatment is preferably in the following range: (1100±25° C. for 0.5 to 5 h)+(1160±25° C. for 1.5 to 10 h)+(1200±25° C. for 2 to 20 h).

The solution heat treatment used was 1165° C. for 7 min. However, this temperature can be adjusted (lower or higher) as well as the holding time with the ultimate goal being to dissolve all secondary phases that may have formed during processing. Thus, the solution heat treatment is preferably expressed as follows: 1165±25° C. for 400±100 sec.

The temperatures and holding times for both steps of the aging heat treatment can be adjusted. The ultimate goal is the form small and blocky M₂₃C₆ carbides uniformly distributed along the grain boundaries during the first step and fine γ′ precipitates during the second step. Several aging heat treatments were employed for testing and consisted of that specified for the commercial alloy (SA): 1010° C. for 2 h followed by air cooling and then 788° C. for 8 h followed by air cooling.

Another commonly used aging heat treatment, the single step aging (SSA), was used to prevent a bimodal size distribution of the γ′ precipitates and consisted of the following:

800° C. for 4 h followed by air cooling. The bimodal size distribution of the γ′ precipitates originated from the higher γ′ solvus temperature of alloy Q compared to the baseline alloy. An alternative dual step aging heat treatment is specified with a first step above the γ′ solvus temperature in the following range: 1060+20° C. for 0.5 to 10 h followed by air cooling and then 800±15° C. for 2 to 15 h followed by air cooling. 

What is claimed:
 1. A nickel alloy, comprising: 18.5 to 20.5% Cr; 9.0 to 11.0% Co; 4.0 to 5.0% Mo; 3.25 to 4.25% W; 2.3 to 2.7% Ti; 1.5 to 1.9% Al; 0 to 0.1% C; 0 to 0.02% B; up to 1.5% Fe, 0.5% or less of Mn, Si, P, Cu or other transition metals; the balance Ni and impurities.
 2. The alloy of claim 1, comprising: 18.5 to 20.5% Cr; 9.0 to 11.0% Co; 4.0 to 5.0% Mo; 3.25 to 4.25% W; 2.3 to 2.7% Ti; 1.5 to 1.9% Al; 0.003 to 0.01% B; 0.0 to 1.5% Fe; 0.04 to 0.08% C; 0 to 0.3% Mn; 0 to 0.15% Si; 0 to 0.15% P; 0 to 0.01% Cu; the balance Ni and impurities.
 3. The alloy of claim 1 comprising M₂₃C₆ carbides at grain boundaries in the alloy.
 4. The alloy of claim 1 comprising γ′ (or γ″ or γ′/γ″) precipitates distributed within a y grain.
 5. The alloy of claim 1 characterizable by a time to failure at 760° C. and 259 MPa of at least 4000 h, or at least 5000 h, or at least 6000 h, or at least 7000 h, or in the range of 4000 to 8000 h or in the range of 4500 to 7500 h.
 6. The alloy of claim 2 characterizable by a time to failure at 800° C. and 155 MPa of at 20 least 3000 h, or at least 4000 h, or at least 5000 h, or at least 6,000 h, or in the range of 3000 to 7000 h or in the range of 4000 to 6500 h.
 7. A method of heat treating a nickel alloy comprising Cr and Co, comprising the sequential steps of treating the nickel alloy: 1100±25° C. for 0.5 to 5 h; 1160±25° C. for 1.5 to 10 h; and 1200±25° C. for 2 to 20 h.
 8. The method of claim 7 wherein the nickel alloy comprises the alloy of claim
 1. 9. The method of claim 7 further comprising a heat treatment comprising heating at 1165±25° C. for 400±100 sec and aging steps of 1010±20° C. for 0.5 to 10 h and then 788±15° C. for 2 to 15 h.
 10. The method of claim 7 further comprising a heat treatment comprising heating at 1165±25° C. for 400±100 sec and aging steps of 1060±20° C. for 0.5 to 10 h and then 800±15° C. for 2 to 15 h.
 11. The method of claim 7 further comprising a heat treatment comprising heating at 1165±25° C. for 400±100 sec and an aging step of 800±15° C. for 2 to 15 h.
 12. The method of claim 7 wherein, relative to HAYNES® 282® alloy, the treatments increase life of the alloy in creep at 760° C. and 259 MPa by at least 70%, or at least 100%, or in the range of 80 to 150%, 90 to 140% or 80 to 130%.
 13. The method of claim 7 wherein the treatments increase the elongation to failure of the alloy in creep at 780° C. and 207 MPa by at least 40%, or at least 60%, or in the range of 40 to 100%, 50 to 90% or 40 to 80%.
 14. A nickel alloy, comprising: 18.5 to 20.5% Cr; 9.0 to 11.0% Co; greater than 50% Ni and characterizable by one or any combination of the following properties: a time to failure at 760° C. and 259 MPa of at least 4000 h, or at least 5000 h, or at least 6000 h, or at least 7000 h, or in the range of 4000 to 8000 h or in the range of 4500 to 7500h; and a time to failure at 800° C. and 155 MPa of at least 3000 h, or at least 4000 h, or at least 5000 h, or at least 6000 h, or in the range of 3000 to 7000 h or in the range of 4000 to 6500 h.
 15. The alloy of claim 14 comprising γ′ (or γ″ or γ′/γ″) precipitates distributed within a y matrix grain interior.
 16. The alloy of claim 14 characterizable by a time to failure at 760° C. and 259 MPa of at least 6000 h.
 17. The alloy of claim 14 characterizable by a time to failure at 800° C. and 155 MPa of at least 4000 h. 5
 18. A nickel alloy, comprising: 18.5 to 20.5% Cr; 9.0 to 11.0% Co; greater than 50% Ni and characterizable by one or any combination of the following properties: a time to failure at 800° C. and 259 MPa of at least 700 h, or at least 800 h, or at least 900 h, or at least 1000 h, or in the range of 700 to 2000 h or in the range of 900 to 1500 h; and a time to failure at 840° C. and 155 MPa of at least 1200 h, or at least 1400 h, or at least 1600 h, or at least 1800 h, or in the range of 1200 to 2200 h or in the range of 1600 to 2000 h.
 19. The alloy of claim 18 comprising γ′ (or γ″ or γ′/γ″) precipitates distributed within a γ matrix grain interior. 